Author’s Accepted Manuscript

Characterization of precipitates in an Al-Zn-Mgalloy processed by ECAP and subsequent annealing

Mohamed A. Afifi, Ying Chun Wang, PedroHenrique R. Pereira, Yangwei Wang, Shukui Li,Yi Huang, Terence G. Langdon

PII: S0921-5093(17)31553-8DOI: https://doi.org/10.1016/j.msea.2017.11.091Reference: MSA35806

To appear in: Materials Science & Engineering A

Received date: 5 September 2017Revised date: 5 November 2017Accepted date: 22 November 2017

Cite this article as: Mohamed A. Afifi, Ying Chun Wang, Pedro Henrique R.Pereira, Yangwei Wang, Shukui Li, Yi Huang and Terence G. Langdon,Characterization of precipitates in an Al-Zn-Mg alloy processed by ECAP andsubsequent annealing, Materials Science & Engineering A,https://doi.org/10.1016/j.msea.2017.11.091

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Characterization of precipitates in an Al-Zn-Mg alloy processed by ECAP

and subsequent annealing

Mohamed A. Afifi

a, Ying Chun Wang

a,b, Pedro Henrique R. Pereira

c, Yangwei Wang

a,b,

Shukui Lia,b

, Yi Huangc, Terence G. Langdon

c

a School of Materials Science and Engineering, Beijing Institute of Technology, Beijing 100081,

China b National Key Laboratory of Science and Technology on Materials under Shock and Impact,

Beijing 100081, China c Materials Research Group, Faculty of Engineering and the Environment,

University of Southampton, Southampton SO17 1BJ, UK

Abstract

Experiments were conducted to examine the influence of equal-channel angular pressing

(ECAP) and post-ECAP annealing on the microstructures of an Al-Zn-Mg alloy. The results

show that precipitates, mainly of the η', η (MgZn2), T (Al20Cu2Mn3) and E (Al18Mg3Cr2) phases,

are fragmented to fine spherical precipitates during ECAP processing for 4 and 8 passes. After

post-ECAP annealing at 393 and 473 K for 20 h, precipitates with larger sizes lie primarily along

the grain boundaries and finer particles are evenly distributed within the grains. Increasing the

numbers of ECAP passes from 4 to 8 leads to an increase in the volume fraction of the finer

precipitates in the ECAP-processed and annealed alloy. After 4 passes and heat treatment at 473

K, the precipitates are slightly larger compared with the alloy processed under identical

conditions and annealed at 393 K. Nevertheless, significant coarsening is evident after

processing for 8 passes and increasing the annealing temperature from 393 to 473 K. Different

types of precipitates are effective in impeding grain growth during the post-ECAP annealing

even at 473 K for 20 h. In addition, η precipitates form within the T and E phases after both

ECAP and post-ECAP annealing.

Keywords: Al-Zn-Mg alloy; Equal-channel angular pressing (ECAP); Grain stability; Heat

treatment; Precipitates.

Corresponding author: Ying Chun Wang, e-mail: wangyc@bit.edu.cn

1

1. Introduction

High strength 7xxx series (Al–Zn–Mg) alloys are widely used in the aerospace and

automobile industries as structural materials [1,2]. Processing by severe plastic deformation

(SPD) as in equal-channel angular pressing (ECAP) [3-6] is recognized as an effective strategy

to enhance the mechanical properties of Al alloys by producing ultrafine-grained (UFG)

microstructures with high densities of dislocations. These dislocations are reorganized during the

multiple passes of ECAP to form fine subgrains or ultrafine grains [6,7]. In addition, second

phase particles may develop in Al alloys during SPD processing [8,9]. These particles consist

primarily of fine precipitates and dispersoid phases in Al-Zn-Mg alloys [10]. During aging

treatment, the precipitation process follows a fixed transformation sequence, given by coherent

GP zones→semi-coherent η'→incoherent η, which is controlled by the precipitation kinetics and

thermodynamics [11]. Dispersoids are intermetallic particles formed usually from elements such

as Cr, Mn and Zr that have low solubility in Al alloys at all temperatures, as with Al3Zr

precipitates [11]. Rod-like η precipitates can be broken into small spherical particles during

ECAP processing and these are thermally stable and very effective in restraining grain growth

during heating [12-14]. Accordingly, the presence of fine precipitates and dispersoids

significantly improves the thermal stability of deformed microstructures [15-21].

Several reports are now available describing the evolution of precipitates and dispersoids

in Al alloys during SPD processing and subsequent heat treatments, and these data provide

information on their underlying properties [8,9,22-25]. For example, GP zones nucleate in the

7050 Al alloy during one pass of ECAP at room temperature and it was shown that the presence

of these ordered phases, allied with the high density of dislocations after ECAP processing,

accelerates the nucleation and growth of η' and η precipitates during post-ECAP aging at 433 K

2

for 60 h [22]. There is also a report that the needle-like precipitates in an AA6060 alloy were

fragmented during ECAP processing up to 8 passes and numerous small precipitates were

dispersedly distributed after ECAP for 2 passes and subsequent aging [23].

In practice, dispersoids have different thermal stabilities and pinning capabilities

according to their sizes and coherency with the Al matrix. For example, the Cr-rich phase in a

7075 Al alloy was thermally stable while the Mg-Zn particles gradually coarsened with

increasing numbers of passes during processing by accumulative roll bonding (ARB) [24].

Another report showed that the high density of Al20Cu2Mn3 precipitates in an Al-Cu-Li alloy

processed by hot rolling was unable to provide sufficient pinning forces to prevent the migration

of grain boundaries during post-deformation heat treatment at 808 K by comparison with the

Al3Zr dispersoids [25]. The lower pinning forces provided by the Mn-rich phase in relation to the

Zr-rich particles is attributed to its relatively lower coherency with the Al matrix, its higher

aspect ratio and its larger size [25]. A review of these data shows that the high strains imposed

by SPD processing are able to increase the precipitation rate and the refinement of different

phases. However, it is not clear whether different phases in the Al-Zn-Mg alloy, including the

MgZn2 (, ), Cr and Mn-rich (E and T) phases, are able to provide a thermally stable structure

during post-ECAP annealing. In addition, the influence of post-ECAP annealing on the

interactions among these different phases is also not clear. It is concluded, therefore, that a more

complete understanding of the precipitation behavior during ECAP and post-ECAP annealing is

needed in order to successfully tailor the precipitation in Al-Zn-Mg alloys and thereby achieve

optimum properties.

The present investigation was initiated specifically to study the formation and distribution

of different types of precipitates in an Al-Zn-Mg alloy during processing by ECAP and after

3

post-ECAP annealing. The effect of these phases on the grain stability was also examined at

temperatures up to 473 K in order to evaluate the potential for applying UFG Al-Zn-Mg alloys in

high temperature environments.

2. Experimental material and procedures

The experiments were conducted using an Al-Zn-Mg alloy bought from Shanghai Silu

Metal Materials Co.Ltd. having a chemical composition, in wt.%, of Al-4.53 Zn-2.52 Mg-0.35

Mn-0.2 Cr-0.11 Cu-0.1 Zr received after solid solution processing at 743 K for 1 h and aging at

393 K for 24 h. Billets of the as-received material with diameters of 9.8 mm and lengths of 65

mm were processed by 4 or 8 ECAP passes using a solid die having an internal channel angle of

90° and an outer arc of curvature of 20°. These angular values impose an equivalent strain of ~1

in each pass [26] and the ECAP processing was performed at 423 K using route Bc where the

billet is rotated by 90⁰ in the same sense between each pass [27]. The ECAP was conducted at

423 K to avoid problems with cracking at lower temperatures [13, 28, 29]. There is experimental

evidence that the temperature of the billet increases during the pressing operation [30-32] and,

using the available data, this temperature rise was estimated as ~18 K for a pressing of 30 s.

Following ECAP processing, samples of the Al-Zn-Mg alloy were annealed at 393 K or

473 K for 20 h using a forced convection furnace and then cooled in air. A pictorial description

of the heat treatment and processing sequence of the Al-Zn-Mg alloy is given in Fig. 1.

Phase identification and microstructural observations were performed using transmission

electron microscopy (TEM) with an FEI Tecnai G2 F20 X-TWIN equipped with a scanning unit

and an energy-dispersive (EDS) detector operating at 200 kV. The samples for TEM were

prepared by electro-polishing in an electrolyte of 8% perchloric acid and 92% ethanol at 248 K

using a voltage of 22 V and a current of ~40 mA. After the electro-polishing, the samples were

4

further thinned by ion-milling (Leica RES 101) at 3 kV using a current of 1 mA for 30 min and

removing any possible contamination caused by electro-polishing. Samples for energy-dispersive

X-ray (EDX) analysis were subjected to additional surface cleaning using Ar/O2 gas chemistry

with a Gatan Plasma cleaning system. Particle size measurements were conducted in scanning

transmission electron microscopy (STEM) mode using a high-angle annular dark field detector

(HAADF). The average particle size was measured by calculating the mean of ~400 particles in

each condition [33]. The average grain sizes of the samples were measured with the aid of TEM

and STEM images by counting more than 150 intercepts using the Circular intercept method

following ASTM E112-12 [34]. In order to quantify the density and volume fraction of

precipitates, the foil thickness measurements were performed with the aid of convergent beam

electron diffraction (CBED). The TEM region was tilted to obtain a two-beam condition with a

dark Kikuchi line crossing the 000Al bright disk and a bright Kikuchi line crossing the 220Al dark

disk in the CBED. A detail description on the thickness determination is given elsewhere [35].

3. Experimental results

3.1 Characterization of the precipitates after ECAP processing

HAADF-STEM images of the as-received and the ECAP-processed Al-Zn-Mg alloy

viewed along the <114>Al zone axis are displayed in Fig. 2. The different phases visible in Fig. 2

were identified through EDX analysis and detailed structural information is shown in Table 1

[36-39]. The average grain size of the as-received Al-Zn-Mg alloy was ~1.3 μm and this was

refined to ~200 nm after 4 passes and ~190 nm after 8 passes of ECAP, respectively [40]. It is

apparent from Fig. 2 (a) that plate and spherical-like second phase particles are heterogeneously

distributed within the microstructure of the as-received alloy. Precipitates mainly of the η' phase

were detected with the presence of a few η, T and E precipitates along the grain boundaries and

5

within the grains of the material in the as-received condition as shown in Fig. 2 (d). Overlapping

of some precipitates was observed which may originate from coarsening during the aging

treatment. The fine η' precipitates were mainly at the grain boundaries and had an average size of

~120 nm whereas the T and E phases were of the order of ~180 nm.

Figure 2 (b) shows that the sizes of the T and E phases are reduced and more η

precipitates form after 4 passes of ECAP by comparison with the unprocessed material. Fig. 2 (e)

is a selected area at higher magnification in Fig. 2 (b) and it shows η, T and E precipitates of ~60

nm as well as numerous fine η' precipitates which are uniformly distributed. The amount of

precipitation further increases by increasing the number of ECAP passes from 4 to 8, as shown in

Fig. 2 (c) and (f). The presence of spherical precipitates of the η, T and E phases having an

average size of ~65 nm together with homogeneous fine η' is apparent in Fig. 2 (f). Furthermore,

the phase was identified within some T and E phases after ECAP in Fig. 2 (e) and (f). Table 2

shows EDX results of selected T and E phases and it reveals a decrease in the content of Mn and

Cr, which are the T and E forming elements, accompanied by an increase of Mg and Zn, which

are the η forming elements, with increasing numbers of ECAP passes. This indicates that the E

and T phases are refined and begin to dissolve to form η phase [41] as is readily visible in Fig. 3

which depicts EDX maps of the Al-Zn-Mg alloy after ECAP for 4 passes showing an E phase

along a grain boundary containing fine η precipitates.

In order to investigate the Al3Zr dispersoid evolution during ECAP processing, Fig. 4

shows TEM and selected area electron diffraction (SAED) images of the Al-Zn-Mg alloy in the

as-received condition and after 4 and 8 passes of ECAP. The SAED patterns were taken along

the <110>Al zone axis of the as-received material and they contain weak spots of Al3Zr at 1/2 of

{220} as shown in Fig. 4 (a) [42]. After 4 and 8 ECAP passes, Al3Zr dispersoid spots are also

6

observed at 1/2 of {220} as displayed in Figs 4 (b) and 4 (c), respectively. This suggests that

there is a small quantity of nano-scaled Al3Zr dispersoids in the microstructure of the Al-Zn-Mg

alloy and these fine precipitates remain stable during the ECAP processing.

3.2 Characterization of the precipitates after post-ECAP annealing

The HAADF-STEM images in Fig 5 display the distributions of second phase particles in

the Al alloy processed by ECAP and further heat treated for 20 h at 393 K or 473 K taken along

the <100>Al zone axes. The bright field TEM along <110>Al in Fig. 6 shows the grain structure

after post-ECAP annealing. It is readily apparent from these images that there are large

precipitates mostly distributed along the grain boundaries after post-ECAP annealing together

with very fine particles evenly distributed within the grains. Inspection of Fig. 5 (a) shows many

fine spherical η' and η precipitates with average sizes of ~13.5 nm in the grains and mainly of T

and E precipitates with diameters of ~82 nm along the grain boundaries after 4 ECAP passes and

further annealing at 393 K for 20 h. Fig. 6 (a) shows that the microstructure of the alloy has fine

equiaxed grains of ~ 230 nm in size. Increasing the annealing temperature to 473 K for the Al-

Zn-Mg alloy processed by 4 ECAP passes produces a slight increase in grain size and a slight

coarsening of the phases by comparison with the material heat treated at 393 K. The size of the

grains in Fig. 6 (b) is ~240 nm and the mean size of the precipitates in the grains is ~21 nm but

this value is ~114 nm along the grain boundaries as depicted in Fig. 5 (b).

After ECAP processing for 8 passes and annealing at 393 K for 20 h, a higher area

fraction of fine precipitates homogeneously distributed within the microstructure was observed

in the micrograph shown in Fig. 5 (c) by comparison with Fig. 5 (a). Spherical precipitates with

average size of ~14 nm were distributed in the grains together with larger grain along the grain

boundaries with size of ~76 nm. Fig. 6 (c) displays fine grains with mean sizes of ~210 nm. The

7

Al-Zn-Mg alloy processed by 8 passes of ECAP and further annealed at 473 K exhibited plate-

like precipitates along the grain boundaries, mostly of the T, E and η phases, with lengths of

~140 nm together with a high fraction of fine η' and η precipitates having an average size of ~ 11

nm uniformly distributed in the Al matrix with a grain size of ~230 nm as shown in Figs 5 (d)

and 6 (d). A comparison of the grain size before and after post-ECAP annealing at 393 and 473

K for 20h reveals small variations which demonstrate that second phase particles are capable of

hindering grain growth during the heat treatments. The volume fraction of the particles after

post-ECAP heat treatments was evaluated following standard procedures [43]. It is 0.085, 0.095

and 0.089 in the microstructure after 4 passes of ECAP followed by annealing at 393 K, at 473

K and after 8 passes of ECAP followed by annealing at 393 K, respectively. For the 8 passes

followed by annealing at 473 K, precipitates are spherical and there are platelet morphologies.

The volume fraction of the spherical precipitates is ~0.072 and ~0.019 for the platelet

precipitates. Table 3 presents a summary of the microstructural features in the as-received

material, after ECAP and after post-ECAP heat treatments.

Figure 7 shows an EDX map in which two η phases having irregular shapes with a mean

size of ~20 nm are located at the edges of a Mn-rich precipitate with a triangular morphology.

The ratio of Zn:Mg detected by EDX was ~2:1 which is in agreement with the composition of

the Laves phase MgZn2 and the Al-Zn-Mg phase diagram [44]. In addition, EDX composition

plots of the Mn and Cr phases and their average compositions are depicted in Fig. 8. These plots

display the percentage of Zn (wt%) detected from the interior of the phases in the post-ECAP

heat-treated Al-Zn-Mg alloy. Each point in Fig. 8 represents the Zn content in a randomly

selected precipitate. By comparing Fig. 8 (a) with Fig. 8 (b), it is apparent that the Zn (wt%)

content in the Al18Mg3Cr2 phase (the E phase) increases with increasing annealing temperature.

8

Furthermore, the results also show that the Cr (wt%) content decreases as a result of the

increasing wt% of Mg and Zn. This may follow from the decomposition of the Al18Mg3Cr2

precipitate (E phase) to form MgZn2 ( phase). Similarly, Fig. 8 (c) and 8 (d)) shows that the Zn

(wt%) in the Al18Cu2Mn3 precipitates increases at the expense of the wt% Mn and Cu which also

indicates that the Al18Cu2Mn3 (T phase) starts to decompose and form MgZn2 ( phase) in the

grain interior. Additionally, it should be noted that the Zn (wt%) was slightly higher after the

post-ECAP heat treatment in the material processed earlier by a larger number of ECAP passes.

Figure 9 shows high magnification TEM micrographs and SAED patterns revealing the

presence of Al3Zr dispersoids within the microstructure of the Al-Zn-Mg alloy processed by

ECAP and further heat treated at 473 K for 20 h. The SAED pattern along the <110> Al zone axis

shows spots of Al3Zr oriented at 1/2 {200} in Fig. 9 (a) and slightly outside 1/2 {220} in Fig. 9

(b). There remain very few Al3Zr dispersoids within the matrix after the post-ECAP heat

treatments. Inspection of Figs 4 and 9, demonstrates that these dispersoids are thermally stable

displaying a spherical shape with a diameter of ~12 nm as marked by black arrows in Fig. 9. In

addition, through observations of the microstructure by TEM, it was found there are twins within

some T and E precipitates as is apparent in Fig. 9 (b).

3.3 Overlapping of the phases

Overlapping of different phases was observed in the Al-Zn-Mg alloy in the as-received

condition and after post-ECAP annealing. Fig. 10 shows a sequence of HRTEM images

revealing the overlapping process in the Al alloy processed by 4 ECAP passes and further heat

treated at 393 K. In Fig. 10 (a), two E phases overlap with a T phase observed close to an array

of dislocations. Fine η precipitates are also evident within the T phase with a size of ~12 nm as

shown in Fig. 10 (b). The highlighted lattice fringes in Fig. 10 (c) show a mismatch of the lattice

9

planes between the two different phases. The FFT patterns displayed in Fig. 10 (d) display the T

and E phase planes with the presence of η phase along (201)Al.

3.4 Twinning of precipitates

Twinning is observed in the material processed by 8 passes of ECAP and further heat

treated at 393 K as depicted in Fig. 11. Fig. 11 (a) shows twinning planes in an E phase which

overlaps with a T phase and Fig. 11 (b) displays clear twinning planes along {112} in the E

phase. The FFT patterns display the extra spots of the twinning plane in addition to the reflection

planes of the coalesced phases. Fig. 11 (c) is a selected IFFT and clearly confirms the presence

of dislocation pairs within the twinning while groups of dislocations are visible in Figs. 11 (d)

and (e). The presence of these dislocations supports the proposal that the twins originate from the

dissociation of dislocations [45].

4. Discussion

4.1 Precipitate evolution during ECAP

In the present investigation, a low number of heterogeneous T and E phases were

detected by EDX in the as-received Al-Zn-Mg alloy. These precipitates have spherical and plate-

like shapes in two-dimensional projections as shown in Fig. 2 (a). Because of the large plastic

straining during ECAP, the precipitates are fragmented, refined and begin to dissolve in the Al

matrix [46]. Accordingly, the presence of fine-scale and homogeneously distributed spherical η'

precipitates after 4 passes of ECAP may result from the occurrence of fragmentation during the

ECAP processing [47]. Increasing the numbers of ECAP passes to 8 leads to the formation of a

larger quantity of spherical η' and η phase [48].

During ECAP processing, the high density of mobile dislocations serves as fast diffusion

paths for solute atoms in the fragmented precipitates [49]. Consequently, the η phase detected

10

within the E and T phases as in Table 2 and Fig. 3 can be attributed to dislocation glide during

straining. In related studies, it was reported that the excess energy from the solute-core

interactions and the relatively open structure of the cores promotes solute diffusion along the

dislocation cores [50,51]. Very few Al3Zr dispersoids were observed in the Al-Zn-Mg alloy in

the as-received and ECAP-processed conditions. Furthermore, the fine scale Al3Zr dispersoids

produced through the aging treatment in the as-received material remain essentially unaffected

after ECAP processing due to a the relatively large solubility limit of Zr in Al of 0.14 wt.%

[52,53].

4.2 Precipitate evolution during post-ECAP annealing

Micrographs of the Al-Zn-Mg alloy after processing by ECAP for 4 passes and heat

treatment at 393 K for 20 h show high amounts of T, E, spherical η and η' phases as is apparent

by inspection of Figs 5 and 6. The presence of very fine precipitates after heating at 393 K may

be ascribed to the precipitation of the dissolved solutes after ECAP processing which forms

newly fine-scale and uniformly distributed particles. Precipitate dissolution during SPD and the

nucleation of precipitates after post-SPD heat treatments were observed in binary Al alloys such

as Al-Cu and Al-Mg [54,55]. Increasing the post-ECAP annealing temperature in the Al alloy

processed by 4 ECAP passes to 473 K leads to a slight coarsening of precipitates. For the

material processed by 8 passes of ECAP, the size and amount of precipitation increases after

heating. The increased number of precipitates is due to the rise of the stored energy after ECAP

processing for 8 passes which leads to an increase of the diffusion rate during the heating process

and subsequently to an increase in the fraction of precipitates [56]. The presence of twin planes

within some precipitates is a direct consequence of the subsequent thermal instability and

coarsening during heating [57, 58]. Fig. 11 displays the dissociation of dislocations to twinning

11

in an E precipitate where groups of dislocations were able to provide cross-slip. The twin planes

are bounded with dislocations along {112} [59]. The presence of dislocations on both twin

boundaries supports the proposal that these dislocations transform into stacking faults and then

twinning is produced with further heating through the coalescence of these faults. The presence

of a dislocation pair within the twinning planes as in Fig. 11 (c) supports the theory of the

multiplication of faults [60]. The higher density of dislocations in the Al-Zn-Mg alloy processed

by 8 passes of ECAP leads to the presence of more twin planes within the precipitates after

heating by comparison with the material processed by 4 ECAP passes as shown in Fig. 9 [40].

It is also revealed in the current study that precipitates positioned along the grain

boundaries are coarser in size and this may lead to some overlapping. During the precipitation

process near grain boundaries, solute atoms diffuse towards the grain boundaries through volume

diffusion, they move along the boundaries to the precipitates and then they are deposited over the

surface of the growing particles through interfacial diffusion [61]. Thus, the growth rates of the

particles at grain boundaries are higher than within the grains and some of these particles have

irregular shapes after overlapping. By contrast, the few Al3Zr dispersoids observed after ECAP

processing were thermally stable after further heat treatments as these dispersoids have an Ll2

structure which is thermally stable up to 698 K [62].

Precipitates of the η phase are evident within the T and E phases after post-ECAP heat

treatments. This may be ascribed to the permanence of η precipitates within these phases from

the ECAP-processed metal or to their nucleation during post-ECAP annealing. The presence of

atomic disorder associated with the grain boundaries contributes to the reduction of the

activation energy barrier for nucleation and this accelerates the precipitation kinetics. Therefore,

grain boundaries act as preferable sites for heterogeneous nucleation of precipitates [63].

12

Accordingly, the T and E phases along the grain boundaries are favorable sites for the

heterogeneous nucleation of η phases during post-ECAP heat treatments as is evident in Fig. 7.

4.3 The effect of precipitates on grain size

The average pinning force applied by the precipitates on the moving boundaries during

grain growth, Z, may be estimated through the relationship [64]

Z = k ( (1)

where k is a scaling factor (1.5 for spherical particles [25]), γ is the energy of the boundary

pinned by the precipitates, 0.32 J m-2

for the alloy [65], r is the mean radius of the precipitates,

and f is the volume fraction of the particles. Substituting of r and f obtained in the results into Eq.

(1) gives Z 0.544 × 106 Nm

-2 , 0.48 × 10

6 Nm

-2 and 0.65 × 10

6 Nm

-2 after ECAP processing for

4 passes and subsequent annealing for 20 h at 393 and 473 K and after 8 ECAP passes followed

by annealing for 20 h at 393K, respectively. For 8 passes and annealing at 473 K, the Z value is

a combination of the spherical and platelet morphologies. For the spherical precipitates, the

pinning force is 0.42 × 106 Nm

-2. The equation for the pinning force by plate precipitates [66]

gives σs ε0.47

where σs is the force exerted by a sphere of the same volume and ε is the aspect ratio

of the particle (length/width). By calculation, the pinning force provided by the plate particles is

0.2 × 106 Nm

-2.It follows from Eq. (1) that the Zener pinning force increases by increasing the

volume fraction (f) and decreasing the particle size (r). There is a small variation in the grain size

before and after heat treatment at 393 and 473 K for 20 h. The coarsening of precipitates along

the grain boundaries in the post-ECAP annealed alloy at 473 K has no significant influence on

the Z values. This is attributed to the increase in the number of fine η' precipitates during heating

at 473 K which sufficiently compensates for the decrease in the pinning force values due to the

increase in size of the remaining phases. Overall, the grain structures of the UFG Al-Zn-Mg

13

alloy display only a slight coarsening after heating up to 473 K for 20 h and the microstructural

stability of the ECAP-processed material is further improved by the presence of fine precipitates

which effectively impede grain boundary migration during any long-term heat treatment.

5. Summary and conclusions

1. During 4 passes of ECAP processing for an Al-Zn-Mg alloy with precipitates of ~120

nm, the η' , η, T and E phases are fragmented into groups of fine spherical particles having

average sizes of ~ 60 nm. Increasing the numbers of passes to 8 increases the numbers of these

precipitates and increases their average size to ~ 65 nm.

2. After post-ECAP annealing at 393 and 473 K for 20 h, the precipitates in the

microstructure consist of larger particles mostly distributed along the grain boundaries and finer

particles distributed evenly within the grains. Increasing the ECAP passes from 4 to 8 leads to

an increase in the quantity of the finer precipitates. Increasing the annealing temperature from

393 to 473 K produces slight coarsening of the precipitates for 4 passes of ECAP but there is

greater coarsening after 8 passes of ECAP.

3. There are a few Al3Zr dispersoids in the as-received material and after ECAP and the

post-ECAP heat treatments. These dispersoids are thermally stable during the heat treatments

and display fine spherical morphology.

4. Due to dislocation motion during ECAP processing, the η phase is nucleated within the

T and E phases where these dislocations act as a fast diffusion path for solute atoms. During

post-ECAP heat treatments, the η phase is also promoted for nucleation in the T and E phases

14

Acknowledgements

This work was supported by the National Natural Science Foundation of China under

Grants nos. 51671030 and 11521062, the Brazilian Federal Agency for the Support and

Evaluation of Graduate Education (CAPES) and the European Research Council under ERC

Grant Agreement no. 267464-SPDMETALS.

15

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Figure captions

Fig. 1. Illustration of the heat treatment and processing sequence of the Al-Zn-Mg alloy.

Fig. 2. HAADF- STEM micrographs of the Al-Zn-Mg alloy in (a,d) the as-received condition

and after ECAP processing for (b,e) 4 and (c,f) 8 passes; (d), (e) and (f) are selected areas at

higher magnifications in (a), (b) and (c), respectively.

Fig. 3. EDX maps of the Al-Zn-Mg alloy after ECAP for 4 passes showing an E phase along a

grain boundary containing fine MgZn2 precipitates.

Fig. 4. TEM micrographs of Al-Zn-Mg alloy in (a) the as-received condition and after ECAP

processing for (b) 4 and (c) 8 passes.

Fig. 5. HAADF-STEM images along <100> Al zone axis showing second phase particles after

ECAP processing for 4 passes and subsequent annealing for 20 h at (a) 393 and (b) 473 K and

after 8 ECAP passes followed by annealing for 20 h at (c) 393 and (d) 473 K.

Fig.6. BF TEM along <110> Al after ECAP processing for 4 passes and subsequent annealing for

20 h at (a) 393 and (b) 473 K and after 8 ECAP passes followed by annealing for 20 h at (c) 393

and (d) 473 K.

Fig. 7. EDX map showing an irregular precipitate with two-η phases along a grain boundary of

the Al-Zn-Mg alloy processed by 8 passes of ECAP and further heat treated at 473 K for 20 h.

Fig. 8. Zn wt% concentration in T (Al18Mg3Cr2) precipitates of the Al-Zn-Mg alloy processed by

(a) 4 and (b) 8 passes of ECAP and further heat treated at 393 and 473 K for 20 h and Zn wt%

concentration in E (Al20Cu2Mn3) phases of the Al-Zn-Mg alloy processed by (c) 4 and (d) 8

passes of ECAP and further heat treated at 393 and 473 K for 20 h.

Fig. 9. TEM images showing precipitates in the Al-Zn-Mg alloy after ECAP processing for (a) 4

and (b) 8 passes followed by heat treatment at 473 K for 20 h.

25

Fig. 10. HRTEM images of the Al-Zn-Mg alloy after ECAP processing for 4 passes and heat

treatment at 393 K for 20 h showing (a) overlapped T and E phases, (b) T phase containing fine

η precipitates, (c) a lattice mismatch between the E and the T phases and (d) FFT pattern of the T

and E phase.

Fig. 11. HRTEM images of twin planes of overlapped T and E phases in (a) and (b) higher

magnification of the twin plane and FFT pattern showing the twinning planes, (c) IFFT showing

dislocation pair, (d) and (e) IFFT presenting a group of dislocations surrounding the twinning

planes.

Table captions

Table 1. Summary of structure information of second phases in Al-Zn-Mg alloys [36-39].

Phase Composition Crystal structure Space group Lattice parameter

η' Zn:Mg = 1:1

to 1.5:1

Monoclinic

Hexagonal P6m2

a = 0.497 nm, c = 0.554 nm

a = 0.496 nm, c = 1.402 nm

η MgZn2 hcp P63/mmc a = 0.516-0.522 nm ,

c = 0.849-0.860 nm

E Al18Mg3Cr2 fcc Fd m a = 1.4526 nm.

T Al20Cu2Mn3 B-centered

orthorhombic structure Bbmm

a = 2.398 nm,

b = 1.254 nm, c = 0.766 nm

Al3Zr Al3Zr Tetragonal I4/mmm a = 4.005 nm,

c = 17.285 nm

Table 2. Element contents of random selected T and E particles in the as-received and ECAP-

processed Al-Zn-Mg alloy detected by EDX (in wt.%)

State phase Al Mg Mn Cr Cu Zn

As-received T 80 1 10.4 0.8 4.1 0.8

E 84 8.2 0.8 4.5 0.05 0.78

ECAP for 4 passes T 83.1 1.8 6.6 1.3 4.1 2.6

E 84.3 9.7 0.3 4.2 1.3 2.2

ECAP for 8 passes T 87 1.9 5.9 0.8 3.1 2.5

E 86 9.5 0.4 3.9 1 1.8

26

Table 3 Summary of microstructural features in the as-received, ECAP-processed and post-ECAP annealed Al-Zn-

Mg alloy

As-

received

ECAP ,

4 passes

ECAP,

8 passes

ECAP, 4

passes +

Annealing,

393 K×20 h

ECAP, 4

passes +

Annealing,

473 K×20 h

ECAP, 8

passes +

Annealing,

393 K×20 h

ECAP, 8

passes +

Annealing,

473 K×20 h

Grain size 1.3 μm 200 nm 190 nm 230 nm 240 nm 210 nm 230 nm

Precipitate

s type

( size/nm)

G.P. (5)

η' (120)

T & E

(180)

G.P.,η',

η, T, E

(~60)

η', η, T,

E

(~65)

η', η

(~13.5)

T, E

(~82)

η',η

(~21)

T, E

(~114)

η', η

(~14)

T , E

(~76)

η' , η

(~11)

η, T, E

(Length:~140)

27

11 1 011

11 1 011

11 1 011

11 1

011

11 1

011

11 1 011

28

.

29

30

31

32

33

34

111